Iron-based rare-earth nanocomposite magnet and method for producing the magnet

ABSTRACT

The iron-based rare-earth nanocomposite magnet of the present invention has a composition T 100−x−y−z−n Q x R y Ti z M n , where T is Fe or a transition metal element in which Fe is partially replaced by Co and/or Ni; Q is B and/or C; R is at least one rare-earth element including substantially no La or Ce; and M is at least one metal element selected from Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb. x, y, z and n satisfy 5≦x≦10 at %, 7≦y≦10 at %, 0.1≦z≦5 at % and 0≦n≦10 at %, respectively. The magnet includes R 2 Fe 14 B-type compound phases and α —Fe phases forming a magnetically coupled nanocomposite magnet structure. The R 2 Fe 14 B-type compound phases have an average crystal grain size of 30 nm to 300 nm and the α —Fe phases have an average crystal grain size of 1 nm to 20 nm. The magnet has magnetic properties including a coercivity of at least 400 kA/m and a remanence of at least 0.9 T.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to an iron-based rare-earth nanocompositemagnet and a method for producing such a magnet. The present inventionalso relates to a rapidly solidified alloy to make an iron-basedrare-earth nanocomposite magnet and to a bonded magnet including aniron-based rare-earth nanocomposite magnet powder.

2. Description of the Related Art

A nanocomposite permanent magnet, having a nanocrystalline structure inwhich a hard magnetic phase such as Nd₂Fe₁₄B₁ phase (which will besometimes referred to herein as a “2-14-1 phase”) and soft magneticphases such as an iron-based boride and α-Fe are magnetically coupledtogether, is currently under development. In the “2-14-1 phase”, Nd maybe replaced with any other rare-earth element, Fe may be partiallyreplaced with Co and/or Ni, and B may also be partially replaced with C(carbon).

The applicant of the present application discovered that when Ti wasadded to an alloy with a particular composition, the nucleation andgrowth of the α —Fe phase could be restricted and the crystal growth ofthe 2-14-1 phase could be advanced preferentially while the molten alloywas being cooled. And the applicant of the present application discloseda nanocomposite magnet, having a structure in which the 2-14-1 phase isdistributed uniformly in the fine iron-based boride and α —Fe phases asa result of the addition of Ti, and a method for producing such a magnetin Patent Document No. 1.

The Ti-containing nanocomposite magnet disclosed in Patent Document No.1, of which the soft magnetic phases are mostly iron-based borides, hasa coercivity of about 500 kA/m to about 1,000 kA/m, which is extremelyhigh for a nanocomposite magnet, but can exhibit a remanence of at mostabout 0.9 T.

Recently, in the field of electronic products including small-sizedmotors and sensors, magnets that have a higher remanence than the magnetdisclosed in Patent Document No. 1 are in high demand. To increase theremanence, the percentage of the α —Fe phase, having a higher saturationflux density than the 2-14-1 phase or the Fe—B phase, may be increased.

Patent Documents Nos. 2 and 3 disclose rare-earth nanocomposite magnetsincluding α —Fe as their main phase. A nanocomposite magnet of that typewould achieve a high remanence exceeding 0.9 T.

-   -   Patent Document No. 1: Japanese Patent No. 3264664    -   Patent Document No. 2: Japanese Patent Application Laid-Open        Publication No. 8-162312    -   Patent Document No. 3: Japanese Patent Application Laid-Open        Publication No. 10-53844

However, the conventional α —Fe based nanocomposite magnets disclosed inPatent Documents Nos. 2 and 3 have such a low coercivity of 400 kA/m orless that it is difficult to actually use such magnets in variousproducts.

SUMMARY OF THE INVENTION

In order to overcome the problems described above, an object of thepresent invention is to provide an iron-based rare-earth nanocompositemagnet exhibiting magnetic properties including a coercivity of at least400 kA/m and a remanence of at least 0.9 T.

Another object of the present invention is to provide a rapidlysolidified alloy to make such an iron-based rare-earth nanocompositemagnet and a powder of the iron-based rare-earth nanocomposite magnet.

An iron-based rare-earth nanocomposite magnet according to the presentinvention has a composition represented by the general formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where T is either Fe alone or atransition metal element in which Fe is partially replaced by at leastone element selected from the group consisting of Co and Ni; Q is atleast one element selected from the group consisting of B and C; R is atleast one rare-earth element including substantially no La or Ce; and Mis at least one metal element selected from the group consisting of Al,Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb. Themole fractions x, y, z and n satisfy the inequalities of 5 at %≦x≦10 at%, 7 at %≦y≦10 at %, 0.1 at %≦z≦5 at % and 0 at %≦n≦10 at %,respectively. The magnet includes R₂Fe₁₄B-type compound phases and α —Fephases that form a magnetically coupled nanocomposite magnet structure.The R₂Fe₁₄B-type compound phases have an average crystal grain size of20 nm or more and the α —Fe phases are present in a grain boundaryregion between the R₂Fe₁₄B-type compound phases and have a thickness of20 nm or less. The magnet has magnetic properties including a coercivityof at least 400 kA/m and a remanence of at least 0.9 T.

In one preferred embodiment, the R₂Fe₁₄B-type compound phases have anaverage crystal grain size of 30 nm to 300 nm and the α —Fe phases havean average crystal grain size of 1 nm to 20 nm.

In another preferred embodiment, the ratio of the average crystal grainsize of the R₂Fe₁₄B-type compound phases to that of the α —Fe phases is2.0 or more.

In still another preferred embodiment, the α —Fe phases are present atgrain boundary triple points of the R₂Fe₁₄B-type compound phases.

In yet another preferred embodiment, the α —Fe phases account for atleast 5 vol % of the overall magnet.

A rapidly solidified alloy to make an iron-based rare-earthnanocomposite magnet according to the present invention has acomposition represented by the general formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where T is either Fe alone or atransition metal element in which Fe is partially replaced by at leastone element selected from the group consisting of Co and Ni; Q is atleast one element selected from the group consisting of B and C; R is atleast one rare-earth element including substantially no La or Ce; and Mis at least one metal element selected from the group consisting of Al,Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb. Themole fractions x, y, z and n satisfy the inequalities of 5 at %≦x≦10 at%, 7 at %≦y≦10 at %, 0.1 at %≦z≦5 at % and 0 at %≦n≦10 at %,respectively. The alloy has an average thickness of 50 μm to 300 μm, andthe alloy includes at least 20 vol % of R₂Fe₁₄B-type compound phaseswith an average crystal grain size of 80 nm or less.

In one preferred embodiment, the alloy has a thickness with a standarddeviation σ of 5 μm or less.

In this particular preferred embodiment, the alloy includes acrystallized layer at least on a free cooling side thereof.

A bonded magnet according to the present invention includes a powder ofthe iron-based rare-earth nanocomposite magnet described above.

A method for producing an iron-based rare-earth nanocomposite magnetaccording to the present invention includes the steps of: preparing amolten alloy having a composition represented by the general formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where T is either Fe alone or atransition metal element in which Fe is partially replaced by at leastone element selected from the group consisting of Co and Ni; Q is atleast one element selected from the group consisting of B and C; R is atleast one rare-earth element including substantially no La or Ce; and Mis at least one metal element selected from the group consisting of Al,Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb, themole fractions x, y, z and n satisfying the inequalities of 5 at %≦x≦10at %, 7 at %≦y≦10 at %, 0.1 at %≦z≦5 at % and 0 at %≦n≦10 at %,respectively; rapidly cooling and solidifying the molten alloy to make arapidly solidified alloy including at least 20 vol % of R₂Fe₁₄B-typecompound phases with an average crystal grain size of 80 nm or less; andheating the rapidly solidified alloy, thereby making an iron-basedrare-earth nanocomposite magnet including the R₂Fe₁₄B-type compoundphases and α —Fe phases that form a magnetically coupled nanocompositemagnet structure, where the R₂Fe₁₄B-type compound phases have an averagecrystal grain size of 20 nm or more, the α —Fe phases are present in agrain boundary region between the R₂Fe₁₄B-type compound phases and havea thickness of 20 nm or less, and the magnet has magnetic propertiesincluding a coercivity of at least 400 kA/m and a remanence of at least0.9 T.

In one preferred embodiment, the R₂Fe₁₄B-type compound phases have anaverage crystal grain size of 30 nm to 300 nm and the α —Fe phases havean average crystal grain size of 1 nm to 20 nm.

In another preferred embodiment, the step of rapidly cooling includesquenching and solidifying the molten alloy to make a rapidly solidifiedalloy with an average thickness of 50 μm to 300 μm and with a thicknessstandard deviation σ of 5 μm or less.

A method of making an iron-based rare-earth nanocomposite magnet powderaccording to the present invention includes the steps of: preparing arapidly solidified alloy according to any of the preferred embodimentsdescribed above to make an iron-based rare-earth nanocomposite magnet;and pulverizing the rapidly solidified alloy into a magnet powder.

In one preferred embodiment, the method further includes the step ofheating the rapidly solidified alloy that has not been pulverized yet orthat has already been pulverized, thereby making a magnet powderincluding R₂Fe₁₄B-type compound phases with an average crystal grainsize of 30 nm to 300 nm and α —Fe phases with an average crystal grainsize of 1 nm to 20 nm and exhibiting magnetic properties including acoercivity of 400 kA/m or more and a remanence of 0.9 T or more.

According to the present invention, by adding Ti, the molten alloy canbe quenched at a lower melt quenching rate than conventional ones withthe nucleation and growth of α —Fe restricted. As a result, a rapidlysolidified alloy including at least 20 vol % of R₂Fe₁₄B-type compoundphases with as small an average crystal grain size as 80 nm or less canbe obtained without amorphizing the entire as-quenched solidified alloy.Even if such a rapidly solidified alloy is heated and crystallized, Tican work so as to restrict the excessive growth of the α —Fe and to growthe R₂Fe₁₄B-type compound phases preferentially. Eventually, the α —Feprecipitates on the grain boundary between the R₂Fe₁₄B-type compoundphases (typically at grain boundary triple points). However, the averagecrystal grain size of the α —Fe phase is 1 nm to 20 nm, which is muchsmaller than that of the R₂Fe₁₄B-type compound phases in the range of 30nm to 300 nm. As a result, magnetic properties including a coercivity ofat least 400 kA/m and a remanence of at least 0.9 T are realized.

BRIEF DESCRIPTION OF DRAWINGS

Portions (a) through (d) of FIG. 1 schematically illustrate thestructures of nanocomposite magnets.

FIGS. 2( a) and 2(b) illustrate a melt-quenching machine (melt spinningmachine in this example).

FIG. 3 is a transmission electron microscope (TEM) photograph showing across-sectional structure of a specific example of the presentinvention.

DETAILED DESCRIPTION OF PREFERRED EMBODIMENTS

According to the present invention, first, a molten alloy having acomposition represented by the general formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n) is prepared. In this formula, T iseither Fe alone or a transition metal element in which Fe is partiallyreplaced by at least one element selected from the group consisting ofCo and Ni; Q is at least one element selected from the group consistingof B and C; R is at least one rare-earth element including substantiallyno La or Ce; and M is at least one metal element selected from the groupconsisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W,Pt, Au and Pb. The mole fractions x, y, z and n satisfy the inequalitiesof 5 at %≦x≦10 at %, 7 at %≦y≦10 at %, 0.1 at %≦z≦5 at % and 0 at %≦n≦10at %, respectively.

Next, the molten alloy with such a composition is rapidly cooled andsolidified to make a rapidly solidified alloy including at least 20 vol% of R₂Fe₁₄B-type compound phases with an average crystal grain size of80 nm or less. This process step will be referred to herein as a “rapidcooling process step”.

In a conventional α —Fe/R₂Fe₁₄B based nanocomposite magnet, the moltenalloy is cooled at as high a quenching rate as possible, thereby makinga substantially amorphous rapidly solidified alloy. Meanwhile, accordingto the present invention, the molten alloy is cooled at a relatively lowquenching rate, thereby making a rapidly solidified alloy including atleast 20 vol % of R₂Fe₁₄B-type compound phases with an average crystalgrain size of 80 nm or less. If the quenching rate were just decreasedwithout adding Ti, the α —Fe would grow faster and earlier than theR₂Fe₁₄B-type compound phase. Thus, in the resultant rapidly solidifiedalloy, the α —Fe phase with a greater average crystal grain size thanthe R₂Fe₁₄B-type compound phase would be its main phase, i.e., thevolume percentage of the α —Fe phase would be higher than that of theR₂Fe₁₄B-type compound phase. Also, the size of the α —Fe phase issignificantly variable with the quenching rate. That is why to producenanocomposite magnets having excellent magnetic properties with goodreproducibility, attempts have been made to form a desired structurewith good hard magnetic properties by making an amorphous rapidlysolidified alloy once and then controlling the heat treatmentconditions.

On the other hand, according to the present invention, not the fullyamorphized rapidly solidified alloy but a rapidly solidified alloy,including at least 20 vol % of R₂Fe₁₄B-type compound phases with anaverage crystal grain size of 80 nm or less, is obtained just after therapid cooling process is finished. The nucleation and the growth of theα —Fe are restricted by adding Ti and the volume percentage of the α —Fephase included in the rapidly solidified alloy is smaller than that ofthe R₂Fe₁₄B-type compound phase included there.

In one preferred embodiment, the thicknesses of the rapidly solidifiedalloys are controlled to have a standard deviation σ of 5 μm or less(i.e., rapidly solidified alloys showing little variation in thicknessare obtained), thereby forming a fine metal structure in which theR₂Fe₁₄B-type compound phase and the amorphous phase are presentuniformly. When the molten alloy is rapidly cooled by the chill roller,the quenching rate achieved by the chill roller varies with thethickness of the rapidly solidified alloy to be obtained. To form auniform nanocomposite magnet structure, the as-quenched, rapidlysolidified alloy should have a uniform quenching rate. For that purpose,the rapid cooling conditions need to be adjusted such that the thicknessof the as-quenched rapidly solidified alloy has a standard deviation σof 5 μm or less. To realize such rapid cooling, a melt spinning processor a melt-quenching process that uses a tundish with multiple tubularholes as disclosed in Japanese Patent Application Laid-Open PublicationNo. 2004-122230 can be used effectively. When any of these rapid coolingprocesses is carried out, a chill roller with a smooth surface ispreferably used and the rapid cooling process is preferably performedwithin a low pressure atmosphere so as to minimize the absorption of theatmospheric gas into the roller.

Thereafter, by thermally treating the rapidly solidified alloy that hasbeen made in this manner, the amorphous phases in the rapidly solidifiedalloy can be crystallized and a nanocomposite magnet with excellentmagnetic properties can be obtained in the end. According to the presentinvention, by heating the rapidly solidified alloy, a nanocompositemagnet structure, including R₂Fe₁₄B-type compound phases with an averagecrystal grain size of 30 nm to 300 nm and α —Fe phases with an averagecrystal grain size of 1 nm to 20 nm, is completed through the heattreatment process. As a result, an iron-based rare-earth nanocompositemagnet, exhibiting magnetic properties including a coercivity of 400kA/m or more and a remanence of 0.9 T or more, can be obtained.

Thus, in the iron-based rare-earth nanocomposite magnet of the presentinvention, the average crystal grain size of the R₂Fe₁₄B-type compoundphases is larger than that of the α —Fe phases. In a preferredembodiment, the α —Fe phase of the very small size is present at thegrain boundary triple point of the R₂Fe₁₄B-type compound phase.

Hereinafter, the difference in structure between the nanocompositemagnet of the present invention and the conventional nanocompositemagnet will be described with reference to FIG. 1.

The rectangular areas on the left-hand side of portions (a) through (d)of FIG. 1 schematically illustrate the nanocrystalline structures ofrapidly solidified alloys, while those on their right-hand sideschematically illustrate the nanocrystalline structures of thermallytreated nanocomposite magnets.

More specifically, portions (a) through (d) of FIG. 1 illustrate ananocomposite magnet according to the present invention, a conventionalα —Fe/R₂Fe₁₄B based nanocomposite magnet, a conventional α —Fe/R₂Fe₁₄Bbased nanocomposite magnet including an additive Ti, and an iron-basedboride/R₂Fe₁₄B based nanocomposite magnet including an additive Ti,respectively.

According to the present invention, the α —Fe cannot be identifiedclearly in the rapidly solidified alloy but a structure in whichR₂Fe₁₄B-type compound phases of a very small size are dispersed inamorphous phases is produced as shown in portion (a) of FIG. 1.Alternatively, α —Fe of a very small size may have been produced at thisstage. This is because by adding Ti, the growth rate of the α —Fe phasewill be restricted during the subsequent heat treatment process, and astructure in which α —Fe phases of a very small size (as indicated bysolid circles in portion (a) of FIG. 1) are present at the grainboundary triple points of the R₂Fe₁₄B-type compound phases that havegrown preferentially will be obtained when the heating process isfinished.

In the conventional α —Fe/R₂Fe₁₄B based nanocomposite magnet, therapidly solidified alloy is almost totally amorphous as shown in portion(b) of FIG. 1. When the heating process is finished, a structure inwhich R₂Fe₁₄B-type compound phases and α —Fe phases of almost the samesize are present as a mixture will be obtained.

It was reported that in the prior art, a metal element such as Ti shouldbe added to the melt of a material alloy to reduce the sizes of theconstituent phases of the nanocomposite magnet structure shown inportion (b) of FIG. 1. When such a metal element is added, a structurein which the sizes of the α —Fe phases (i.e., the black portions inportion (c) of FIG. 1) and the R₂Fe₁₄B-type compound phases (i.e., thewhite portions in portion (c) of FIG. 1) have been both reduced can beobtained as shown in portion (c) of FIG. 1 by controlling the conditionsof the heating process.

In the iron-based boride/R₂Fe₁₄B based nanocomposite magnet including anadditive Ti, not the α —Fe/R₂Fe₁₄B based nanocomposite magnet describedabove, R₂Fe₁₄B-type compound phases are produced in the rapidlysolidified alloy as shown in portion (d) of FIG. 1. However, in theresultant structure, those R₂Fe₁₄B-type compound phases are surroundedby iron-based boride (Fe—B) phases thinly.

In the conventional α —Fe/R₂Fe₁₄B based nanocomposite magnet, the sizeof the α —Fe phases cannot be made smaller than that of the R₂Fe₁₄B-typecompound phases. Thus, in the known α —Fe/R₂Fe₁₄B based nanocompositemagnet, the α —Fe phases and R₂Fe₁₄B-type compound phases havesubstantially the same average crystal grain size. Also, to improve theperformance of the nanocomposite magnet, the size of the α —Fe phasesneeds to be reduced. Thus, a nanocomposite magnet, in which the sizes ofthe α —Fe phases (i.e., the black portions in portion (c) of FIG. 1) andthe R₂Fe₁₄B-type compound phases (i.e., the white portions in portion(c) of FIG. 1) were both reduced uniformly by controlling the conditionsof the rapid solidification process and heat treatment process tocrystallize the material alloy, was reported as shown in portion (c) ofFIG. 1.

Unlike these conventional nanocomposite magnets, the nanocompositemagnet of the present invention has a distinct structure in which α —Fephases of a very small size are dispersed discretely on the grainboundary between R₂Fe₁₄B-type compound phases of a relatively largesize, thus realizing excellent nanocomposite magnet performance. Such astructure is realized by not just adding Ti to the alloy but alsodispersing the R₂Fe₁₄B-type compound phases of a small sizeintentionally in the as-quenched rapidly solidified alloy. As a result,the R₂Fe₁₄B-type compound phases will grow faster and earlier than the α—Fe phases and the α —Fe phases of a very small size will not benucleated at the grain boundary triple points until the R₂Fe₁₄B-typecompound phases have grown sufficiently.

In the iron-based boride/R₂Fe₁₄B based nanocomposite magnet shown inportion (d) of FIG. 1, the iron-based boride (Fe—B) phases are presentin the form of a film. On the other hand, according to the presentinvention, the α —Fe phases of a very small size are dispersed at thegrain boundary triple points as shown in portion (a) of FIG. 1. Thereason why the α —Fe phases are dispersed in this manner is not clearyet.

Hereinafter, a preferred embodiment of a method for producing aniron-based rare-earth nanocomposite magnet according to the presentinvention will be described.

First, a configuration for a melt-quenching machine for use in thispreferred embodiment will be described with reference to FIG. 2.

[Melt Quenching Machine]

In this preferred embodiment, a material alloy is prepared by using amelt quenching machine such as that shown in FIG. 2. The alloypreparation process is performed within an inert atmosphere to preventthe material alloy, which includes rare-earth element R and Fe that areeasily oxidizable, from being oxidized. The inert gas may be either arare gas of helium or argon, for example, or nitrogen.

The machine shown in FIG. 2 includes material alloy melting andquenching chambers 1 and 2, in which a vacuum or an inert atmosphere ismaintained at an adjustable pressure. Specifically, FIG. 2( a)illustrates an overall arrangement of the machine, while FIG. 2( b)illustrates a portion of the machine on a larger scale.

As shown in FIG. 2( a), the melting chamber 1 includes: a melt crucible3 to melt, at an elevated temperature, a material 20 that has been mixedto have a desired magnet alloy composition; a reservoir 4 with a teemingnozzle 5 at the bottom; and a mixed material feeder 8 to supply themixed material into the melt crucible 3 while maintaining an airtightcondition. The reservoir 4 stores the melt 21 of the material alloytherein and is provided with a heater (not shown) to maintain thetemperature of the melt teemed therefrom at a predetermined level.

The quenching chamber 2 includes a rotating chill roller 7 for quenchingand solidifying the melt 21 that has been dripped through the teemingnozzle 5.

In this machine, the atmosphere and pressure inside the melting andquenching chambers 1 and 2 are controllable within prescribed ranges.For that purpose, atmospheric gas inlet ports 1 b, 2 b and 8 b andoutlet ports 1 a, 2 a and 8 a are provided at appropriate positions ofthe machine. In particular, the gas outlet port 2 a is connected to apump to control the absolute pressure inside the quenching chamber 2within a range of 30 kPa to the normal pressure (i.e., atmosphericpressure).

The melt crucible 3 may define a desired tilt angle to pour the melt 21through a funnel 6 into the reservoir 4. The melt 21 is heated in thereservoir 4 by the heater (not shown).

The teeming nozzle 5 of the reservoir 4 is positioned on the boundarywall between the melting and quenching chambers 1 and 2 to drip the melt21 in the reservoir 4 onto the surface of the chill roller 7, which islocated under the nozzle 5. The orifice diameter of the teeming nozzle 5may be 0.5 mm to 2.0 mm, for example. If the viscosity of the melt 21 ishigh, then the melt 21 cannot flow through the teeming nozzle 5 easily.In this embodiment, however, the pressure inside the quenching chamber 2is kept lower than the pressure inside the melting chamber 1.Accordingly, an appropriate pressure difference is created between themelting and quenching chambers 1 and 2, and the melt 21 can be teemedsmoothly.

To achieve a good thermal conductivity, the chill roller 7 may be madeof Al alloy Cu alloy, carbon steel, brass, W, Mo or bronze. However, theroller 7 is preferably made of Cu, Fe or an alloy including Cu or Fe,because such a material realizes a sufficient mechanical strength at areasonable cost. Also, if the chill roller is made of a material otherthan Cu or Fe, the resultant rapidly solidified alloy cannot come offthe chill roller easily and might be wound around the roller. The chillroller 7 may have a diameter of 300 mm to 500 mm, for instance. Thewater-cooling capability of a water cooler provided inside the chillroller 7 is calculated and adjusted based on the latent heat ofsolidification and the volume of the melt teemed per unit time.

The machine shown in FIG. 2 can rapidly solidify 10 kg of material alloyin 10 to 20 minutes, for example. The rapidly solidified alloy obtainedin this manner is in the form of an alloy thin strip (or alloy ribbon)22 with a thickness of 10 μm to 300 μm and a width of 2 mm to 3 mm, forexample.

[Melt Quenching Process]

First, the melt 21 of the material alloy, which is represented by thefollowing general formula, is prepared and stored in the reservoir 4 ofthe melting chamber 1 shown in FIG. 2( a).

The alloy has a composition represented by the general formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where T is either Fe alone or atransition metal element in which Fe is partially replaced by at leastone element selected from the group consisting of Co and Ni; Q is atleast one element selected from the group consisting of B and C; R is atleast one rare-earth element including substantially no La or Ce; and Mis at least one metal element selected from the group consisting of Al,Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb. Andthe mole fractions x, y, z and n satisfy the inequalities of 5 at %≦x≦10at %, 7 at %≦y≦10 at %, 0.1 at %≦z≦5 at % and 0 at %≦n≦10 at %,respectively.

Next, the melt 21 is dripped through the teeming nozzle 5 onto the chillroller 7 to contact with, and be quenched and solidified by, the chillroller 7 within a low-pressure Ar atmosphere. In this case, anappropriate rapid solidification technique, making the quenching ratecontrollable precisely, should be adopted.

In this preferred embodiment, the melt 21 is preferably quenched andsolidified at a quenching rate of 1×10⁴° C./s to 1×10⁶° C./s, morepreferably 3×10⁴° C./s to 1×10⁶° C./s, and even more preferably 1×10⁵°C./s to 1×10⁶° C./s.

A period of time during which the molten alloy 21 is quenched by thechill roller 7 is equivalent to an interval between a point in time thealloy contacts with the outer circumference of the rotating chill roller7 and a point in time the alloy leaves the roller 7. In this period oftime, the alloy has its temperature decreased to be a supercooledliquid. Thereafter, the supercooled alloy leaves the chill roller 7 andtravels within the inert atmosphere. While the thin-strip alloy istraveling, the alloy has its heat dissipated into the atmospheric gas.As a result, the temperature of the alloy further drops. In thisembodiment, the pressure of the atmospheric gas is 30 kPa to theatmospheric pressure. Thus, the heat of the alloy can be dissipated intothe atmospheric gas even more effectively, and the Nd₂Fe₁₄B compound cannucleate and grow finely and uniformly in the alloy. It should be notedthat unless an appropriate amount Ti has been added to the materialalloy, then the α —Fe phase nucleates and grows faster and earlier inthe rapidly solidified alloy that has gone through the quenching processdescribed above, thus growing the α —Fe excessively and deterioratingthe magnetic properties of the resultant magnet.

In this preferred embodiment, the surface velocity of the roller isadjusted to fall within the range of 10 m/s to 30 m/s (more preferably14 m/s to 2.5 m/s and even more preferably, 18 m/s to 22 m/s) and thepressure of the atmospheric gas is set to 30 kPa or more to increase thesecondary cooling effects caused by the atmospheric gas. In this manner,a rapidly solidified alloy, including at least 20 volume % of R₂Fe₁₄Btype compound phases with an average crystal grain size of as small asabout 80 nm or less, is obtained. Such a crystallized layer is formedsubstantially uniformly on the free cooling side of the rapidlysolidified alloy thin strip. In addition, another crystallized layer maybe formed thinly on the uppermost surface of the rapidly solidifiedalloy thin strip that has been in contact with the surface of the chillroller. Meanwhile, the intermediate portion between these twocrystallized layers is in an amorphous or quasi-amorphous state.

[Heat Treatment]

In this preferred embodiment, the rapidly solidified alloy is thermallytreated within an argon atmosphere. Preferably, the alloy is heated at atemperature rise rate of 5° C./s to 20° C./s, maintained at atemperature of 550° C. to 850° C. for 30 seconds to 20 minutes, and thencooled to room temperature. This heat treatment results in nucleationand/or crystal growth of metastable phases in a remaining amorphousphase, thus forming a nanocomposite crystalline structure. According tothe present invention, the nanocrystalline Nd₂Fe₁₄B phase alreadyaccounts for at least 20 volume % of the as-cast alloy that has juststarted being thermally treated. Thus, α —Fe and other crystallinephases will not increase their sizes too much and the respectiveconstituent phases other than the nanocrystalline Nd₂Fe₁₄B phase (i.e.,soft magnetic phases) will be dispersed finely and uniformly. Theresultant nanocomposite magnet that has gone through this heat treatmenthas a nanocrystalline structure in which the grain boundary between theNd₂Fe₁₄B crystalline phases is mostly composed of α —Fe, which shouldaccount for at least 5 vol % of the entire magnet. Thus, the magnet as awhole can exhibit increased remanence.

If the heat treatment temperature were lower than 550° C., then a lot ofamorphous phases might remain even after the heat treatment and theresultant coercivity could not reach the desired level depending on theconditions of the quenching process. On the other hand, if the heattreatment temperature exceeded 850° C., the grain growth of therespective constituent phases would advance too much, thus decreasingthe remanence B_(r) and deteriorating the loop squareness of thedemagnetization curve. For these reasons, the heat treatment temperatureis preferably 550° C. to 850° C., more preferably 570° C. to 820° C.

To prevent the alloy from being oxidized, the heat treatment ispreferably conducted within an inert atmosphere. The heat treatment mayalso be performed within a vacuum of 0.1 kPa or less.

The rapidly solidified alloy yet to be heat-treated may includemetastable phases such as Fe₃B, Fe₂₃B₆ and R₂Fe₂₃B₃ phases in additionto the R₂Fe₁₄B-type compound and amorphous phases. In that case, whenthe heat treatment is finished, the R₂Fe₂₃B₃ phase will havedisappeared. Instead, crystal grains of an iron-based boride (e.g.,Fe₂₃B₆), exhibiting a saturation magnetization that is equal to, or evenhigher than, that of the R₂Fe₁₄B phase, or α —Fe phase can be grown.

After the heat treatment, the R₂Fe₁₄B-type compound phases need to havean average crystal grain size of less than 300 nm, which is a singlemagnetic domain size. The R₂Fe₁₄B-type compound phases preferably havean average crystal grain size of 30 nm to 150 nm, more preferably 30 nmto 100 nm, to increase the coercivity and improve the loop squareness ofthe demagnetization curve. On the other hand, if the α —Fe phases had anaverage crystal grain size of more than 20 nm, then the exchangeinteractions among the respective constituent phases would weaken and α—Fe particles with a multiple magnetic domain structure, not the singledomain structure, would produce in increasing numbers, thusdeteriorating the loop squareness of the demagnetization curve anddecreasing magnetization, B_(r), and (BH)_(max). Nevertheless, if theaverage crystal grain size of these phases were less than 1 nm, then ahigh coercivity could not be achieved anymore. In view of theseconsiderations, the α —Fe phases preferably have an average crystalgrain size of 1 nm to 20 nm.

As can be seen, according to the present invention, the average crystalgrain size of the R₂Fe₁₄B-type compound phases is greater than that ofthe α —Fe phases and the ratio of the former to the latter is 1.5 ormore. This ratio is preferably 2.0 or more.

It should be noted that the thin strip of the rapidly solidified alloycould be coarsely cut or coarsely pulverized before subjected to theheat treatment. When the heat treatment is finished, the resultantmagnet is finely pulverized to make a magnet powder. Then, various typesof bonded magnets can be made from this magnet powder by performingknown process steps on the powder. In making a bonded magnet, the magnetpowder of the iron-based rare-earth alloy is mixed with an epoxy ornylon resin binder and then molded into a desired shape. In this case, amagnet powder of any other type (e.g., an Sm—Fe—N based magnet powder orhard ferrite magnet powder) may be mixed with the nanocomposite magnetpowder.

[Why this Composition is Preferred]

As described above, the composition of an iron-based rare-earthnanocomposite magnet according to the present invention is representedby the formula: T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where T is eitherFe alone or a transition metal element including Fe and at least oneelement selected from the group consisting of Co and Ni; Q is at leastone element selected from the group consisting of B and C; R is at leastone rare-earth element including substantially no La or Ce; and M is atleast one metal element selected from the group consisting of Al, Si, V,Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb. And themole fractions x, y, z and n satisfy the inequalities of 5 at %≦x≦10 at%, 7 at %≦y≦10 at % (preferably 8 at %≦y≦10 at %), 0.1 at %≦z≦5 at %(preferably 0.5 at %≦z≦4 at %) and 0 at %≦n≦10 at %, respectively.

Q is either B (boron) only or a combination of B and C (carbon). Theatomic percentage ratio of C to Q is preferably 0.5 or less.

With no Ti added, if the mole fraction x of Q were less than 7 at %,then the amorphous phases would be produced at a much lower rate. As aresult, no uniform nanocrystalline metal structure could be produced anda remanence B_(r) of 0.9 T or more could not be achieved. According tothe present invention, by adding Ti, the amorphous phases can beproduced at an increased rate, and therefore, the mole fraction x of Qhas a lower limit of 5 at %. On the other hand, if the mole fraction xof Q exceeded 10 at %, then the percentage of the α —Fe phase, which hasa higher saturation magnetization than any other constituent phase,would decrease and soft magnetic phases such as Fe₃B, Fe_(3.5)B andFe₂₃B₆ would nucleate and a remanence B_(r) of 0.9 T or more could notbe achieved. In view of these considerations, the mole fraction x of Qis preferably set to fall within the range of 5 at % to 10 at %, morepreferably 5.5 at % to 9.5 at %, and even more preferably 5.5 at % to9.0 at %. The upper limit of a further preferred range of the molefraction x is 8 at %. A portion (up to 50 at %) of B may be replacedwith carbon (C) because the magnetic properties and the metal structurewill not be affected in that case.

R is at least one element selected from the group consisting of therare-earth elements (including Y). Preferably, R includes substantiallyno La or Ce, because the presence of La or Ce would decrease thecoercivity and the loop squareness of the demagnetization curve.However, there will be no problem of degrading the magnetic propertiesif very small amounts (i.e., 0.5 at % or less) of La and Ce are includedas inevitable impurities. Therefore, if the content of La or Ce is 0.5at % or less, then the magnet may be regarded as including substantiallyno La or Ce.

More particularly, R preferably includes Pr or Nd as an indispensableelement, a portion of which may be replaced with Dy and/or Tb. If themole fraction y of R were less than 7 at %, then compound phases havingthe nanocrystalline R₂Fe₁₄B-type structure, which contribute toexpressing coercivity, would not crystallize sufficiently and acoercivity H_(cJ) of 400 kA/m or more could not be realized. On theother hand, if the mole fraction y of R exceeded 10 at %, then thepercentages of the iron-based borides and α —Fe with ferromagneticproperties would both decrease. For these reasons, the mole fraction yof the rare-earth element R is preferably 7 at % to 10 at % (e.g., 7.5at % to 9.8 at %), more preferably 8 at % to 9.8 at %, and mostpreferably 8.2 at % to 9.7 at %.

To achieve the above-described effects, Ti is an indispensable element.The additive Ti increases the coercivity H_(cJ), remanence B_(r) andmaximum energy product (BH)_(max) and improves the loop squareness ofthe demagnetization curve.

If the mole fraction z of Ti were less than 0.1 at %, then the aboveeffects would not be achieved fully even though Ti is added.Nevertheless, if the mole fraction z of Ti exceeded 5 at %, then thevolume percentage of the amorphous phases, remaining even in the alloythat has been heated and crystallized, would increase so much as todecrease the remanence B_(r) easily. In view of these considerations,the mole fraction z of Ti is preferably 0.1 at % to 5 at %. The lowerlimit of a more preferable z range is 0.5 at % and the upper limitthereof is 4 at %. The lower limit of an even more preferable z range is1 at %.

The balance of the magnet, other than the elements described above, maybe a transition metal element T, which consists mostly of Fe.Alternatively, one or two transition metal elements T, selected from thegroup consisting of Co and Ni, may be substituted for a portion of Fe,because the desired hard magnetic properties are achievable in thatcase, too. However, if more than 50% of Fe were replaced with Co and/orNi, then a high remanence B_(r) of 0.5 T or more could not be realized.For that reason, the percentage of Fe replaced is preferably from 0% to50%. Also, by substituting Co for a portion of Fe, the loop squarenessof the demagnetization curve improves and the Curie temperature of theR₂Fe₁₄B phase increases, thus improving the thermal resistance.Furthermore, by adding Co, the molten alloy being quenched has adecreased viscosity, and therefore, the melt quenching process can beperformed with good stability. The percentage of Fe that is replaceablewith Co is preferably 0.5% to 15%.

Optionally, one, two or more additives M, selected from the groupconsisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W,Pt, Au and Pb, may be added. By adding these elements, not onlyimprovement of magnetic properties but also expansion of the best heattreatment temperature range will be achieved as well. However, if themole fraction of M exceeded 10 at %, then the magnetization woulddecrease. For that reason, the mole fraction n of M should fall withinthe range of 0 at % to 10 at %, more preferably 0.1 at % to 5 at %.

EXAMPLES Example 1

For each of Samples Nos. 1 through 21 having the compositions shown inthe following Table 1, Nd, Pr, B, C, Ti, Cu, Ga, Co, Zr, V, Nb and Fewere weighed so that the mixture had a total weight of 30 g and then themixture was put into a nozzle of transparent quartz having an orificewith a diameter of 0.8 mmφ at the bottom.

Thereafter, the materiel in the nozzle was melted by an inductiveheating process within an Ar atmosphere. When the temperature of theresultant molten alloy reached 1,400° C., the molten alloy in the nozzlewas pressurized with an argon gas at 30 kPa, thereby ejecting the moltenalloy through the orifice at the bottom of the nozzle onto the surfaceof the chill roller.

The chill roller was rotated at a high velocity while being cooledinside so that the outer circumference thereof would have itstemperature maintained around at room temperature. Accordingly, themolten alloy, which had been ejected through the orifice, contacted withthe surface of the roller to have its heat dissipated therefrom whilebeing forced to rapidly move in the surface velocity direction. Thechill roller had a surface velocity Vs of 20 m/s.

In this manner, a thin strip of rapidly solidified alloy with an averagethickness of 40 μm to 50 μm and a width of 0.9 mm to 2.0 mm was made. Asa result of an analysis using a powder X-ray diffraction (powder XRD)analyzer, it was confirmed that the rapidly solidified alloy thusobtained had a quenched alloy structure in which amorphous phases andcrystalline phases that would be Nd₂Fe₁₄B and α —Fe phases coexisted. Itwas also confirmed that the thickness of the resultant rapidlysolidified alloy thin strip had a standard deviation a of σ μm or lesseverywhere.

Next, the alloy thin strip thus obtained was cut into a number offlakes, each having a length of about 20 mm, which were then thermallytreated and crystallized by maintaining its temperature at 630° C. to700° C. for 10 minutes.

When the crystalline phases of the thermally treated rapidly solidifiedalloy thin strip were analyzed by a powder XRD, it was discovered thatSamples Nos. 1 to 21 had a metal structure consisting essentially ofNd₂Fe₁₄B and α —Fe phases (but some of them included Fe—B phases such asFe₃B and Fe₂₃B₆ phases). Also, when the fine metal structure wasobserved with a transmission electron microscope, a nanocomposite magnetstructure, including Nd₂Fe₁₄B phases with an average crystal grain sizeof 30 nm to 100 nm and α —Fe phases with as small an average crystalgrain size as 1 nm to 20 nm in the grain boundary region (with athickness of several nm to 20 nm) between the Nd₂Fe₁₄B crystal grains(particularly in the vicinity of the grain boundary triple points), wasidentified. FIG. 3 is a transmission electron microscope (TEM)photograph showing a cross section of Sample No. 19. As can be seen fromthis photograph, the α —Fe phase accounted for at least 5 vol % of theoverall magnet. The magnetic properties of the thermally treated andcrystallized thin strip of the rapidly solidified alloy were measured atroom temperature using a vibrating sample magnetometer and are shown inTable 2.

It should be noted that: if the surface velocity Vs of the chill rollerwas changed within the range of 14 m/s to 18 m/s during the quenchingprocess, a rapidly solidified alloy thin strip with a thickness of 52 μmto 74 μm was obtained. Thus, it was confirmed that when the surfacevelocity Vs of the chill roller decreased, the resultant rapidlysolidified alloy thin strip had an increased thickness but its standarddeviation σ still fell within the range of 2.2 to 4.2.

Comparative Examples

30 g of material, in which Nd, B, Nb, Cr and Fe were mixed together soas to have one of the alloy compositions Nos. 22 to 30 shown in Table 1,was put into a nozzle of transparent quartz having an orifice with adiameter of 0.8 mmφ at the bottom. Thereafter, the material in thenozzle was melted by an inductive heating process within an Aratmosphere. When the temperature of the resultant molten alloy reached1,400° C., the molten alloy in the nozzle was pressurized at 30 kPa,thereby ejecting the molten alloy through the orifice at the bottom ofthe nozzle onto the surface of the chill roller rotating at a velocityVs of 20 m/s. In this manner, the molten alloy was quenched to make arapidly solidified alloy thin strip with a width of 0.9 mm to 1.1 mm andan average thickness of 40 μm to 50 μm.

Next, the alloy thin strip thus obtained was cut into a number offlakes, each having a length of about 20 mm, which were then thermallytreated and crystallized by maintaining its temperature at 630° C. to700° C. for 10 minutes.

TABLE 1 Alloy composition (at %) R T Q Nd Pr Fe Co B C Ti M EXAMPLES 110 0 bal. 0 6 1 3 0 2 9 0 bal. 8 5.5 0.5 1 0 3 9 0 bal. 8 6 0 1 0 4 8.250 bal. 8 5.5 0.5 1 0 5 9 0 bal. 0 5.5 0.5 1 Cu 0.25 6 9 0 bal. 0 5.5 0.51 Ga 0.5 7 9 0 bal. 0 5.5 0.5 1 Nb 0.5 8 9 0 bal. 8 5.5 0.5 0.5 V 0.5 99 0 bal. 8 5.5 0.5 0.5 Zr 0.5 10 9 0 bal. 8 5 0.5 0.5 Nb 0.5 11 0 9 bal.8 5.5 0.5 2 0 12 4.12 4.13 bal. 8 5.5 0.5 3 0 13 0 8.25 bal. 8 5.5 0.5 10 14 7 0 bal. 0 7 0 1 0 15 7.5 0 bal. 0 9 0 3 0 16 8.5 0 bal. 0 7.5 0 10 17 9.8 0 bal. 0 7.5 0 2 0 18 9 0 bal. 3 8 0 2 0 19 9 0 bal. 0 5.5 0.51 0 20 9 0 bal. 8 6 0 0.5 0 21 9 0 bal. 8 6 0 0.3 0 COMP EXAM 22 9 0bal. 0 5 0 0 0 23 8 0 bal. 0 5 0 0 0 24 7 0 bal. 0 5 0 0 0 25 8 0 bal. 06 0 0 0 26 8 0 bal. 0 9 0 0 0 27 7 0 bal. 0 7 0 0 0 28 8 0 bal. 0 10 0 00 29 7 0 bal. 0 7 0 0 Nb 1 30 7 0 bal. 0 7 0 0 Cr 1

When the crystalline phases of the thermally treated rapidly solidifiedalloy thin strip were analyzed by a powder XRD, it was discovered thatSamples Nos. 22 to 30 had a metal structure consisting essentially ofNd₂Fe₁₄B and α —Fe phases. Also, when the fine metal structure wasobserved with a transmission electron microscope, a fine metalstructure, including Nd₂Fe₁₄B phases with an average crystal grain sizeof nm to 100 nm and α —Fe phases, was identified. The magneticproperties of the thermally treated and crystallized thin strip of therapidly solidified alloy were measured at room temperature using avibrating sample magnetometer and are also shown in Table 2.

TABLE 2 Magnetic properties B_(r) H_(cJ) (BH)_(max) (mT) (kA/m) (kJ/m³)EXAMPLES 1 864 871 118 2 1,035 593 149 3 1,013 579 144 4 1,043 502 147 5986 574 138 6 973 595 137 7 983 630 141 8 995 576 137 9 982 551 140 101,023 569 147 11 997 613 146 12 1,045 515 145 13 1,060 541 151 14 1,076420 130 15 910 434 115 16 980 457 120 17 950 783 132 18 950 740 135 19992 581 141 20 970 569 118 21 945 534 111 COMP 22 902 418 74 EXAM 23 960397 84 24 946 265 58 25 913 379 79 26 917 405 97 27 1,035 353 99 28 874417 91 29 1,026 409 117 30 1,046 395 112

As is clear from Tables 1 and 2, excellent magnetic properties,including a coercivity of 400 kA/m or more and a remanence of 0.9 T ormore, were realized in the specific examples of the present invention,whereas the remanence was lower than 0.9 T in the comparative examples.

According to the present invention, an iron-based rare-earthnanocomposite magnet, exhibiting magnetic properties including acoercivity of 400 kA/m or more and a remanence of T or more, is providedand can be used effectively in small-sized motors, sensors and otherelectronic devices that need high-remanence magnets.

1. An iron-based rare-earth nanocomposite magnet having a compositionrepresented by the formula:T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n), where Tis either Fe alone or Fe in combination with at least one elementselected from the group consisting of Co and Ni; Q is at least oneelement selected from the group consisting of B and C; R is at least onerare-earth element including substantially no La or Ce; and M is atleast one metal element selected from the group consisting of Al, Si, V,Cr, Mn, Cu, Zn, Ga, Zr, Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb, the molefractions x, y, z and n satisfying the inequalities of5 at % ≦x ≦8 at %,7 at % ≦y ≦10 at %,0.1 at % ≦z ≦5 at % and0 at % ≦n ≦10 at %, respectively, wherein the magnet includes R₂T₁₄Qcompound phases and α —Fe phases that form a magnetically couplednanocomposite magnet structure, and wherein the R₂T₁₄Q compound phaseshave an average crystal grain size of 20 nm or more and the α —Fe phasesare present at grain boundary triple points in a grain boundary regionbetween the R₂T₁₄Q compound phases, the grain boundary region having athickness of 20 nm or less, wherein a ratio of the average crystal grainsize of the R₂T₁₄Q compound phases relative to that of the α —Fe phasesis 2.0 or more, and wherein the magnet has magnetic properties includinga coercivity of at least 400 kA/m and a remanence of at least 0.9 T. 2.The iron-based rare-earth nanocomposite magnet of claim 1, wherein theR₂T₁₄Q compound phases have an average crystal grain size of 30 nm to300 nm and the α —Fe phases have an average crystal grain size of 1 nmto 20 nm.
 3. The iron-based rare-earth nanocomposite magnet of claim 1,wherein the α —Fe phases account for at least 5 vol % of the overallmagnet.
 4. A bonded magnet including a powder of the iron-basedrare-earth nanocomposite magnet of claim
 1. 5. A method for producing aniron-based rare-earth nanocomposite magnet, the method comprising thesteps of: preparing a molten alloy having a composition represented bythe formula: T_(100−x−y−z−n)Q_(x)R_(y)Ti_(z)M_(n) , where T is either Fealone or Fe in combination with at least one element selected from thegroup consisting of Co and Ni; Q is at least one element selected fromthe group consisting of B and C; R is at least one rare-earth elementincluding substantially no La or Ce; and M is at least one metal elementselected from the group consisting of Al, Si, V, Cr, Mn, Cu, Zn, Ga, Zr,Nb, Mo, Ag, Hf, Ta, W, Pt, Au and Pb, the mole fractions x, y, z and nsatisfying the inequalities of5 at %≦x≦8 at %,7 at %≦y≦10 at %,0.1 at %≦z≦5 at % and0 at %≦n≦10 at %, respectively; rapidly cooling and solidifying themolten alloy to make a rapidly solidified alloy including at least 20vol % of R₂T₁₄ Q compound phases with an average crystal grain size of80 nm or less; and heating the rapidly solidified alloy, thereby makingan iron-based rare-earth nanocomposite magnet including the R₂T₁₄Qcompound phases and α —Fe phases that form a magnetically couplednanocomposite magnet structure, where the R₂T₁₄Q compound phases have anaverage crystal grain size of 20 nm or more, the α —Fe phases arepresent at grain boundary triple points in a grain boundary regionbetween the R₂T₁₄Q compound phases, the grain boundary region having athickness of 20 nm or less, wherein a ratio of the average crystal grainsize of the R₂T₁₄Q compound phases relative to that of the α —Fe phasesis 2.0 or more, and the magnet has magnetic properties including acoercivity of at least 400 kA/m and a remanence of at least 0.9 T. 6.The method of claim 5, wherein the R₂T₁₄Q compound phases have anaverage crystal grain size of 30 nm to 300 nm and the α —Fe phases havean average crystal grain size of 1 nm to 20 nm.
 7. The method of claim5, wherein the step of rapidly cooling includes quenching andsolidifying the molten alloy to make a rapidly solidified alloy with anaverage thickness of 50 μm to 300 μm and with a thickness standarddeviation σ of 5 μm or less.